Nanoscale fiber films, composites, and methods for alignment of nanoscale fibers by mechanical stretching

ABSTRACT

Articles including nanoscale fibers aligned by mechanical stretching are provided. Methods for making composite materials comprising a network of aligned nanoscale fibers are also provided. The network of nanoscale fibers may be substantially devoid of a liquid, and may be a buckypaper. The network of nanoscale fibers also may be associated with a supporting medium.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application is a divisional application of U.S. patent applicationSer. No. 12/690,558, filed Jan. 20, 2010, which claims priority to U.S.Provisional Patent Application No. 61/145,849, filed Jan. 20, 2009,which are incorporated herein by reference.

STATEMENT REGARDING FEDERALLY SPONSORED RESEARCH OR DEVELOPMENT

This invention was made with U.S. government support under Contract No.FA9550-05-1-0271 awarded by the Air Force Office of Scientific Researchand Contract No. N00014-08-M-0348 awarded by the Office of NavelResearch STTR Program. The U.S. government has certain rights in theinvention.

BACKGROUND OF THE INVENTION

This invention relates generally to nanoscale fibers, and moreparticularly to methods for aligning carbon nanotubes or other nanoscalefibers to a high degree of alignment in the production of buckypapercomposite materials.

Carbon nanotubes and nanofibers have both rigidity and strengthproperties, such as high elasticity, large elastic strains, and fracturestrain sustaining capabilities. Such a combination of properties isgenerally not present in conventional materials. In addition, carbonnanotubes and nanofibers are some of the strongest fibers currentlyknown. For example, the Young's Modulus of single-walled carbonnanotubes can be about 1 TPa, which is about five times greater thanthat for steel (about 200 GPa), yet the density of the carbon nanotubesis about 1.2 g/cm³ to about 1.4 g/cm³. The tensile strength ofsingle-walled carbon nanotubes is generally in the range of about 50 GPato about 200 GPa. This tensile strength indicates that compositematerials made of carbon nanotubes and/or nanofibers could likely belighter and stronger as compared to current high-performance carbonfiber-based composites.

In addition to their exceptional mechanical properties, carbon nanotubesand nanofibers may provide either metallic or semiconductorcharacteristics based on the chiral structure of fullerene. Some carbonnanotubes and nanofibers also possess superior thermal and electricalproperties such as thermal stability up to about 2800° C. in a vacuumand about 750° C. in air, thermal conductivity about twice as much asthat of diamond, and an electric current carrying capacity about 1000times greater than that of copper wire. Therefore, carbon nanotubes andnanofibers are regarded as one of the most promising reinforcementmaterials for the next generation of high-performance structural andmultifunctional composites.

Thin films or sheets of nanoscale fiber networks, or buckypapers (BP),offer a promising platform to fabricate high-performance nanoscale fibercomposites because BPs are easy to handle during fabrication of thecomposite, and thus, may be incorporated into conventional compositesprocessing to fabricate nanocomposites.

Nanoscale fibers have both exceptional mechanical and functionalproperties, which conventional macroscopic carbon fibers do not offer.However, four main factors tend to affect the performance ofnanocomposites: 1) nanoscale fiber dispersion, 2) nanoscale fiberalignment, 3) interface bonding between the nanoscale fibers and thecomposite matrix, and 4) aspect ratio of the nanoscale fibers. Forinstance, the composite nanoscale fiber loading may be too low (lessthan 20 wt %), there may be a lack of adequate nanoscale fiberalignment, or the smaller aspect ratios of nanoscale fibers such as CNTs(less than 10,000) may result in poor load transfer between the matrixand CNTs when the composites are under loads.

Methods for aligning nanoscale fibers such as carbon nanotubes includemagnetic field-induced alignment, mechanical stretching of synthesizednanotube forests, shear force-induced alignment, AC electric fieldalignment, electrospinning, and electrophoretic alignments duringnanotube composite fabrication. However, the loose and weakly bondedstructures of nanotube networks make it difficult to uniformly transferforce throughout nanotube networks, thus hindering the development ofpractical methods to further improve nanoscale fiber alignment in BPthrough mechanical stretching to achieve a high degree of alignment(e.g., greater than 20%).

It would therefore be desirable to provide improved nanotube alignmenttechniques for alignment of nanoscale fibers in BP.

SUMMARY OF THE INVENTION

A method for aligning carbon nanotubes or other nanoscale fibers isprovided. In one aspect, the method comprises providing a network ofnanoscale fibers substantially devoid of a liquid and mechanicallystretching the network of nanoscale fibers in a first direction. In oneembodiment, the network is a buckypaper. In certain embodiments, themethod further comprises providing a supporting medium on or in thenetwork of nanoscale fibers before the step of mechanically stretching.The step of mechanically stretching stretches the network of nanoscalefibers and the supporting medium in a first direction.

In some embodiments, the nanoscale fibers comprise carbon nanotubes. Inone embodiment, the carbon nanotubes have an average length of at least1 millimeter.

In another aspect, a method for aligning carbon nanotubes or othernanoscale fibers comprises providing a network of nanoscale fibers,providing a supporting medium on or in the network of nanoscale fibers,and mechanically stretching the network of nanoscale fibers and thesupporting medium in a first direction. In one embodiment, the networkis substantially devoid of a liquid.

In certain embodiments, the supporting medium comprises a flexiblethermoplastic material. In one embodiment, the flexible thermoplasticmaterial comprises a polyethylene film.

In particular embodiments, the method further comprises, after the stepof stretching, removing the supporting medium from the network ofnanoscale fibers. In some embodiments, the step of removing comprisesthermally decomposing the supporting medium. In one embodiment, themethod further comprises annealing the supporting medium before the stepof thermally decomposing.

In yet another aspect, a method for making a composite material isprovided. The method comprises providing a network of nanoscale fibers,mechanically stretching the network of nanoscale fibers in a firstdirection to form a network of aligned nanoscale fibers, andincorporating the network of aligned nanoscale fibers onto or into amatrix material.

In another aspect, an article comprising a network of nanoscale fibersis provided. The network has been mechanically stretched to align atleast a portion of the nanoscale fibers and the network has a Young'smodulus ranging from 5 GPa to 25 GPa, in the direction of the nanoscalefiber alignment. In particular embodiments, the network has a tensilestrength ranging from 200 MPa to 668 MPa in the direction of thenanoscale fiber alignment.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 illustrates an embodiment of an unstretched and stretchedbuckypaper labeled with notations for calculating stretching ratioaccording to Equation 1.

FIG. 2 illustrates an embodiment of a method for mechanically stretchingnanotube buckypaper in SEM micrographs and in an analogous schematicrepresentation.

FIG. 3 shows the stretching setup and stretched BP/PE film used inExample 1.

FIG. 4 is a graph showing the degree of alignment via Raman spectrumcomparison of the stretched (30% stretch ratio) and unstretched BPsamples of Example 1 using a polarized Raman spectrometer (InVia,Renishaw, Inc).

FIG. 5 shows typical stress-strain curves for embodiments of amechanically stretched buckypaper.

FIG. 6 is a graph showing the tensile modulus of embodiments of amechanically stretched buckypapers at various stretch ratios made inExample 1.

FIG. 7 is a graph showing the strength of embodiments of themechanically stretched buckypapers at various stretch ratios made inExample 1.

FIG. 8 shows the SEM images of buckypapers with different stretchingratios.

FIG. 9 shows typical stress-strain curves for the control (unstretched)and stretched composites for epoxy/BP composites made in Example 1.

FIG. 10 shows stress-strain curves for embodiments of mechanicallystretched and unstretched buckypaper made in Example 2.

FIG. 11A shows tensile stress-strain curves for CNTs/BMI composites madein Example 2. FIG. 11B shows the tensile strength and Young's modulus ofthese samples.

FIG. 12 is a graph which shows the comparison of mechanical propertiesbetween the CNTs/BMI composites made in Example 2 with state-of-the-artunidirectional (UD) carbon fiber composites for aerospace structuralapplications.

FIG. 13A depicts the fracture surface morphology of one embodiment of acomposite material at a particular magnification.

FIG. 13B depicts the fracture surface morphology of one embodiment of acomposite material at a particular magnification.

FIG. 13C depicts the fracture surface morphology of one embodiment of acomposite material at a particular magnification.

FIG. 13D depicts the fracture surface morphology of one embodiment of acomposite material at a particular magnification.

FIG. 14A is a graph comparing storage modulus of the CNT_(S)/BMIcomposites made in Example 2. FIG. 14B is a graph showing T_(g)determined by Tan Delta for the BP composites made in Example 2.

FIG. 15 is a graph comparing the electrical conductivity values of theCNTs/BMI composites made in Example 2.

FIG. 16 is a graph showing the best fitting curve of the Raman intensityversus orientation angle as described in Example 3.

FIG. 17 is a graph showing the stress-strain curves of neat MWNT sheetsof different stretch ratios made in Example 3.

FIG. 18 is a graph showing the electrical conductivity measurements ofthe neat MWNT sheets made in Example 3 paralleled to the alignmentdirection.

FIGS. 19A-C are SEM micrographs showing the fracture surface morphologya 40%-stretched specimen after tensile tests as described in Example 3.

FIGS. 20-22 are graphs showing typical tensile stress-strain curves offunctionalized CNT sheet/BMI composites made in Example 4.

FIG. 23 is a graph showing the mechanical properties of CNT sheet/BMIcomposites made in Example 4 as compared to UD carbon fiber reinforcedcomposites.

FIG. 24 is a graph of ATR-FTIR spectra of pristine CNTs, functionalizedCNTs, and pristine and functionalized aligned (40% stretch) CNTsheet/BMI composites made in Example 4 (Trace a: pristine CNT; Trace b:expoxidation functionalized CNT; Trace c: pristine 40% aligned CNT/BMInanocomposite; Trace d: functionalized 40% aligned CNT/BMInanocomposite).

FIG. 25 shows the proposed reaction mechanism for functionalization ofCNTs as described in Example 4.

FIG. 26 is a graph of the Raman spectrometer data for the curingmechanism described in Example 4.

FIG. 27 shows the typical stress-strain curves of CNT sheets reinforcedBMI nanocomposites made in Example 4 along the nanotube alignmentdirection.

FIGS. 28A-B are SEM micrographs showing the fracture surface morphologyof a functionalized 40% stretch alignment specimen after tensile testingas described in Example 4.

FIGS. 29A-B are graphs showing dynamic mechanical analysis (DMA) resultsfor the samples made in Example 4.

FIG. 30 shows the reaction mechanism between the functionalized CNTs andepoxy resin matrix as described in Example 5.

FIG. 31 is a graph showing the curves of degree of functionalization(DOF) versus functionalization time and m-CPBA concentration for thecomposites made in Example 5.

FIG. 32A is a graph showing the attenuated total reflection Fouriertransform infrared (ATR-FTIR) spectrum comparison of the composites madein Example 5. FIG. 32B is a graph of the Raman spectrometer data for thereaction mechanism described in Example 5.

FIGS. 33A (pristine double-walled nanotube) and 34B (functionalizeddouble-walled nanotube) are HRTEM micrographs of the samples made inExample 5.

FIG. 34 is a graph of showing the mechanical properties of resultantnanocomposites of Example 5.

FIG. 35A is a graph showing the load transfer efficiency factor η_(B)logistic fitting to determine η_(B) as described in Example 5; FIG. 35Bis a graph showing the relationship of load transfer efficiency and DOF.

FIGS. 36A (random) and 36B (aligned) are SEM micrographs showing thecross-section of random and aligned CNT sheets made in Example 5.

FIG. 37 is a graph showing the typical stress-strain curves of CNT sheetreinforced epoxy nanocomposites with/without alignment andfunctionalization as made in Example 5.

FIG. 38A depicts a graph of the tensile strength of one embodiment of arandom CNT sheet nanocomposite made by the procedure of Example 5.

FIG. 38B depicts a graph of the Young's modulus of one embodiment of arandom CNT sheet nanocomposite made by the procedure of Example 5.

FIGS. 39A-B are SEM micrographs of the fracture surface morphology of apristine aligned CNT sheet reinforced epoxy composite specimen as madein Example 5. FIGS. 39C-D are SEM micrographs of the fracture morphologyof a functionalized aligned CNT sheet reinforced epoxy composite as madein Example 5. FIG. 39E is the HRTEM image of cross-section of a pristinealigned CNT sheet reinforced epoxy composite as made in Example 5.

FIG. 40A is a graph of the tensile strength of the functionalized andaligned CNT composites made in Example 5 in comparison tostate-of-the-art high-strength unidirectional structural CFRP systems;FIG. 40B is a graph showing the failure strain of the functionalized andaligned CNT composites made in Example 5 in comparison tostate-of-the-art high-strength unidirectional structural CFRP systems.

DESCRIPTION OF THE INVENTION

Methods have been developed for inducing a selected orientation ofnanoscale fibers by mechanical stretching of buckypapers and/ornanoscale fiber composites that include a thermoplastic film.Buckypapers comprising aligned nanoscale fibers, and compositescomprising the same, are also provided. In one aspect, the methods use asupporting medium, such as a flexible polymer film, to protect andsupport nanotube networks in a buckypaper during a mechanical stretchingprocess to further increase nanotube alignment. The supporting mediumeffectively transfers the stretching load to the buckypaper nanotubenetwork uniformly so that the BP does not break under high stretchingstrains (e.g., stretching strains for achieving a high degree ofalignment). As used herein, “degree of alignment” refers to thepercentage of nanoscale fibers in the network which is aligned in aparticular direction of alignment. The degree of alignment may becalculated using Raman spectrometer, X-ray defraction, or any othermethods known in the art.

In certain embodiments, the methods may mechanically stretch abuckypaper without using a supporting medium. No liquid treatment (e.g.,solvent or surfactant treatment) of the buckypaper or network isnecessary to achieve alignment of the nanoscale fibers. These methodsrely on relatively simple mechanical force, requiring inexpensivematerials and equipment to achieve a high degree of nanoscale fiberalignment in the buckypaper, which can lead to inexpensive futurescaling-up and commercialization.

Since nanotubes are highly anisotropic in nature, the alignment ofnanotubes in buckypaper is desirable for achieving strong mechanicalproperties and high electrical and thermal conductivity. In addition,alignment of nanoscale fibers, such as nanotubes, in buckypaper isdesirable for utilizing the exceptional mechanical properties ofnanotubes along their axial direction. Furthermore, enhanced contactefficiency between the nanotubes allows for realization of moreself-assembly and high density packing of the nanotubes, therebyproviding improved load transfer in the composites.

Aligned buckypaper can be impregnated by thermosetting and/orthermoplastic resins, such as epoxy, for fabricating nanocomposites. Useof the aligned BP in composites facilitates realizing the full potentialof carbon nanotubes or nanofiber buckypaper for high mechanicalperformance and multifunctional applications, such as lightweighthigh-performance structural materials, electromagnetic interferenceshielding materials, and thermal management materials, and otherapplications.

Methods of Nanotube Alignment and Composite Production

In some embodiments, the method for aligning nanoscale fibers includesproviding a network of nanoscale fibers and mechanically stretching thenetwork of nanoscale fibers in a first direction. The network ofnanoscale fibers is substantially devoid of a liquid. As seen herein,“substantially devoid of a liquid” means the network comprises liquid inan amount less than about 10 wt % of the network, typically less thanabout 1 wt %, 0.1 wt %, or 0.01 wt %.

Without being bound by a particular theory, the frictional forces orsimilar forces between nanoscale fibers in the network films orbuckypapers that are substantially devoid of a liquid act to distributethe mechanical stretching force more uniformly throughout the network.Thus, the films and buckypapers remain intact at high stretchingstrains. The absence of a liquid or a lubricant would reduce theeffectiveness of the load transfer due to the frictional forces betweenthe nanoscale fibers.

In one embodiment, the network of nanoscale fibers is devoid of solventor surfactants during the step of mechanically stretching.

As used herein, “mechanically stretching” or “mechanically stretch”refers to treatment of sheets of networks of fibers or buckypapers bypulling or applying mechanical loads to the sheets in opposed or offsetdirections. For example, the sheets could be mechanically stretched bypassing the sheets between rollers set a different speeds such that oneroller set at a relatively slower rotational speed “holds” the sheet andthe other side set at a relatively faster rotational speed “stretches”the sheet. In certain embodiments, a multiple stage roll stretchingmachine may be used to for a continuous process of mechanicallystretching buckypaper.

In one embodiment, a buckypaper is stretched using a Shimadzu machine.In such embodiments, the stretching ratio (or stretch ratio, A %) of BPsamples was calculated by Equation 1.

$\begin{matrix}{{\Delta\;\%} = {\frac{L_{2 -}L_{1}}{L_{1} - L_{a} - L_{b}} \times 100{\%.}}} & {{Equation}\mspace{14mu} 1}\end{matrix}$Where L₁, and L₂ are length of a BP strip before and after stretching,L_(a) and L_(b) are the lengths of the segments held by the stretchingclamp as shown in FIG. 1. It should be understood that Equation 1 can beused as stated, or modified to suit the buckypaper shape and theparticular process used to stretch the buckypaper.

In certain embodiments, a buckypaper and supporting media are stretchedtogether to align the nanoscale fibers of the buckypaper. In oneembodiment, the flexible thermoplastic films (i.e., plastic films) areused as supporting media sandwiching a buckypaper to form athermoplastic/nanotube composite film with the aid of heat and pressure.For example, the supporting media and BP sandwich may be compressedusing in a vacuum bag and heated to form a composite film. Then, thecomposite film is mechanically stretched. FIG. 2 illustrates anembodiment of a method for mechanically stretching nanotube buckypaperin SEM micrographs and in an analogous schematic representation.

In some embodiments, the composite film may be heated prior to and/orduring the stretching step, to enhance stretching without breakage.

During stretching, the supporting media hold and bring the nanotubescloser together in the nanotube network to induce nanotube alignmentalong the stretching direction. The supporting media uniformly transferthe stretching force to the nanotubes in the buckypaper to realizealignment and avoid nanotube network breakdown.

In one embodiment, the composite film may be burned off (e.g., in afurnace) after being stretched. As used herein, “burned” or “burning”refers to the thermal decomposition, melting, or degradation of thesupporting media. In other embodiments, a solvent may be used todissolve and wash out the supporting medium (e.g., a polymer supportingmedium).

In particular embodiments, the BP/supporting media composite may beannealed to prevent the composite from shrinking at high temperatures(e.g., during the burning off of the supporting media). Shrinking of thecomposite would undesirably affect the alignment of the stretched BP. Inone embodiment, a BP/polyethylene (PE) film supporting media compositeis annealed in a vacuum bag at 90-100° C. for 10 minutes. Without beingbound by a particular theory, the annealing disorients the PE molecules,which were previously oriented during the stretching process, to preventshrinking of the PE film at high temperatures.

Using a supporting medium for the BP stretching process providesparticular advantages. First, the supporting medium protects and holdsthe nanotube network together during stretching, as interactions amongnanotubes in the BP may be weak, such that the nanotube network maybreak down during stretching and cannot transfer the stretching loaduniformly throughout entire nanotube networks. In addition, thesupporting medium may provide a force to hold the nanotube network toavoid the BP breaking down due to local nanotube slippage duringstretching.

When a polymeric supporting medium is used, the polymer may be infusedinto the nanotube network so that the loading can be transferred moreuniformly to the nanotube network and result in an even deformation ofthe BP during the stretching process. Therefore, in sandwich composites,supporting medium provide a large deformation capability and a uniformtransfer of stretching load to the entire nanotube network to inducenanotube alignment.

Some supporting medium, such as PE, also do not strongly interact andadhere with nanotubes due to the non-reactive natures of both molecules.

In addition, supporting media used in some embodiments may have highbreaking strain (e.g., more than several hundred percent). For instance,a large stretching ratio (>20%) may be realized in a PE/BP compositefilm to achieve a high degree of nanotube alignment.

In addition, certain supporting media may be easily thermally decomposedby heating, leaving almost no residue on the buckypaper, thus resultingin thermally stable materials. Therefore, purification of the stretchedBP materials may be achieved by simple thermal burning processes.

In one embodiment, a method is provided that includes using PE polymeras a supporting medium in a PE/BP composite, mechanically stretching thePE/BP composite, and then burning the PE/BP composite to remove the PEpolymer from the stretched BP/PE composite film to obtain purified BPwith improved nanotube alignment.

In certain embodiments, the stretching deformation is irreversiblebecause the original nanoscale fiber entanglements have been changed dueto the stretching force, and the nanoscale fibers self-assemble into analigned assembly due to van der Waals interaction. Hence, in particularembodiments there is no retraction of the BP after stretching.

Embodiments of the methods have an excellent potential for use in themass production of high-performance nanotube and nanofiber-reinforcedpolymer composites. In some embodiments, the aligned buckypapers undergocomposite fabrication processes for making final composites, such asvacuum-assisted resin transfer molding (VARTM), resin transfer molding(RTM), vacuum infusion process (VIP), autoclave/prepreg process,carbon-carbon impregnation, or a combination thereof.

In addition to using mechanically stretched buckypaper to providealigned nanoscale fibers in composites, the methods described hereinalso provide high loading and larger aspect ratios. Without being boundby a particular theory, the effect of nanoscale fiber orientation on themechanical properties of BP composites is significant because theorientation of the nanotubes allows for control of the contactefficiency, and thus the efficiency of load transfer, between therelatively rigid neighboring nanotubes. Specifically, the mechanicallystretched buckypaper provides nanoscale fibers in a more compacted form,as well being comprised of longer networks of nanoscale fibers which arecapable of being stretched mechanically. In certain embodiments, thenanotube rope sizes are increased and the packing density is higher ascompared to pre-stretched nanotube sheets. These features are furtherenhanced, for example, by using large-aspect-ratio, MWNTs having anaverage length of at least 1 millimeter (“millimeter-long MWNTs”).

In certain embodiments, the method of making a composite furthercomprises impregnating the mechanically stretched buckypaper with aresin and then B-stage curing the resin to form a prepreg. By using (1)a prepregging process to achieve high nanotube loading and goodresin/nanotube impregnation in each thin prepreg layer (10-20 mm)through resin B-stage compression; and (2) a precise control of thenanotube concentration, higher loading of the nanoscale fibers in theresultant composite is realized. In addition, nanoscale fiber filmprepregs may be much easier to handle than untreated buckypaper. Thus,in one embodiment, the method making a composite includes layup of aplurality of aligned buckypaper prepregs and then curing the B-stagecured resin of the prepregs to form a composite.

In certain embodiments, the method further comprises functionalizing thenanoscale fibers. In one embodiment, nanoscale fibers are functionalizedby covalently attaching chemical groups to the nanoscale fiber tofacilitate covalent bonding between the nanoscale fiber and a resinmatrix. Without being bound by a particular theory, covalentfunctionalization creates defects in nanotube lattice, which also lowerselectrical and thermal conductivity of the CNTs. Hence, covalentfunctionalization is a double-edged sword for realizing high mechanicalproperties of CNT reinforced composites. Thus, the degree offunctionalization should balance the increase in the interfacial bondingwith decrease in mechanical properties of CNTs to maximize themechanical properties in the resultant composites. Specific degrees offunctionalization (DOF) that can improve interfacial bonding withoutunduly sacrificing the intrinsic mechanical properties of CNTs.

In certain embodiments, the method for making a composite furthercomprises functionalizing the nanoscale fibers by contacting thenanoscale fibers with a peroxyacid (e.g.,meta-chloromethaneperoxylbenzoic acid or m-chloroperoxybenzoic acid(m-CPBA)) to graft an epoxide group onto the nanoscale fibers.

Nanoscale Fiber Films

As used herein, the term “nanoscale fibers” refers to a thin, greatlyelongated solid material, typically having a cross-section or diameterof less than 500 nm. In certain embodiments, the nanoscale fibers aresingle-walled carbon nanotubes (SWNTs), multiple-walled carbon nanotubes(MWNTs), carbon nanofibers (CNFs), or mixtures thereof. Carbon nanotubesand carbon nanofibers have high surface areas (e.g., about 1,300 m²/g),which results in high conductivity and high multiple internalreflection. In a preferred embodiment, the nanoscale fibers comprise orconsist of carbon nanotubes, including both SWNTs and MWNT. SWNTstypically have small diameters (˜1-5 nm) and large aspect ratios, whileMWNTs typically have large diameters (˜5-200 nm) and small aspectratios. CNFs are filamentous fibers resembling whiskers of multiplegraphite sheets or MWNTs.

In certain embodiments, the nanoscale fibers comprise carbon nanotubesan average length of at least 1 millimeter (available from NanocompTechnologies, Concord, N.H.).

As used herein, the terms “carbon nanotube” and the shorthand “nanotube”refer to carbon fullerene, a synthetic graphite, which typically has amolecular weight between about 840 and greater than 10 milliongrams/mole. Carbon nanotubes are commercially available, for example,from Unidym Inc. (Houston, Tex. USA) or Carbon Nanotechnologies, Inc.(Houston Tex. USA), or can be made using techniques known in the art.

The nanotubes optionally may be opened or chopped, for example, asdescribed in U.S. Pat. No. 7,641,829 B2.

The nanotube and nanofibers optionally may be chemically modified orcoated with other materials to provide additional functions for thefilms produced. For example, in some embodiments, the carbon nanotubesand CNFs may be coated with metallic materials to enhance theirconductivity.

As used herein, the term “nanoscale film” refers to thin, preformedsheets of well-controlled and dispersed porous networks of SWNTs, MWNTs,CNFs, or mixtures thereof. Films of carbon nanotubes and nanofibers, orbuckypapers, are a potentially important material platform for manyapplications. Typically, the films are thin, preformed sheets ofwell-controlled and dispersed porous networks of SWNTs, MWNTs, carbonnanofibers CNFs, or mixtures thereof. The carbon nanotube and nanofiberfilm materials are flexible, light weight, and have mechanical,conductivity, and corrosion resistance properties desirable for numerousapplications. The film form also makes nanoscale materials and theirproperties transferable to a macroscale material for ease of handling.

The nanoscale fiber films be made by essentially any suitable processknown in the art. In one embodiment, the buckypaper is made bystretching or pushing synthesized nanotube “forests” to form sheets orstrips. In another embodiment, the buckypaper is made by consolidationof syntheses nanotube aerogel to form film membranes.

In some embodiments, the nanoscale fiber film materials are made by amethod that includes the steps of (1) suspending SWNTs, MWNTs, and/orCNF in a liquid, and then (2) removing a portion of the liquid to formthe film material. In one embodiment, all or a substantial portion ofthe liquid is removed. As seen herein, “a substantial portion” meansmore than 50%, typically more than 70, 80%, 90%, or 99% of the liquid.The step of removing the liquid may include a filtration process,vaporizing the liquid, or a combination thereof. For example, the liquidremoval process may include, but is not limited to, evaporation (ambienttemperature and pressure), drying, lyophilization, heating to vaporize,or using a vacuum.

The liquid includes a non-solvent, and optionally may include asurfactant (such as Triton X-100, Fisher Scientific Company, NJ) toenhance dispersion and suspension stabilization. As used herein, theterm “non-solvent” refers to liquid media that essentially arenon-reactive with the nanotubes and in which the nanotubes are virtuallyinsoluble. Examples of suitable non-solvent liquid media include water,and volatile organic liquids, such as acetone, ethanol, methanol,n-hexane, benzene, dimethyl formamide, chloroform, methylene chloride,acetone, or various oils. Low-boiling point liquids are typicallypreferred so that the liquid can be easily and quickly removed from thematrix material. In addition, low viscosity liquids can be used to formdense conducting networks in the nanoscale fiber films.

For example, the films may be made by dispersing nanotubes in water or anon-solvent to form suspensions and then filtering the suspensions toform the film materials. In one embodiment, the nanoscale fibers aredispersed in a low viscosity medium such as water or a low viscositynon-solvent to make a suspension and then the suspension is filtered toform dense conducting networks in thin films of SWNT, MWNT, CNF or theirmixtures. Other suitable methods for producing nanoscale fiber filmmaterials are disclosed in U.S. patent application Ser. No. 10/726,074,entitled “System and Method for Preparing Nanotube-based Composites;”U.S. Patent Application Publication No. 2008/0280115, entitled “Methodfor Fabricating Macroscale Films Comprising Multiple-Walled Nanotubes;”and U.S. Pat. No. 7,459,121 to Liang et al., which are incorporatedherein by reference.

Additional examples of suitable methods for producing nanoscale fiberfilm materials are described in S. Wang, Z. Liang, B. Wang, and C.Zhang, “High-Strength and Multifunctional Macroscopic Fabric ofSingle-Walled Carbon Nanotubes,” Advanced Materials, 19, 1257-61 (2007);Z. Wang, Z. Liang, B. Wang, C. Zhang and L. Kramer, “Processing andProperty Investigation of Single-Walled Carbon Nanotube (SWNT)Buckypaper/Epoxy Resin Matrix Nanocomposites,” Composite, Part A:Applied Science and Manufacturing, Vol. 35 (10), 1119-233 (2004); and S.Wang, Z. Liang, G. Pham, Y. Park, B. Wang, C. Zhang, L. Kramer, and P.Funchess, “Controlled Nanostructure and High Loading of Single-WalledCarbon Nanotubes Reinforced Polycarbonate Composite,” Nanotechnology,Vol. 18, 095708 (2007).

In certain embodiments, the nanoscale fiber films are commerciallyavailable nanoscale fiber films. For example, the nanoscale fiber filmsmay be preformed nanotube sheets made by depositing synthesizednanotubes into thin sheets (e.g., nanotube sheets from NanocompTechnologies Inc., Concord, N.H.). MWNT sheets from Nancomp havesubstantial nanotube entanglements and possible interconnection throughNanocomp's proprietary floating catalyst synthesis and aerogel condensemethod. Theses MWNT sheets can reach up to a meter long and arecommercially available, which makes them practical for manufacturingbulk composites.

In various embodiments, good dispersion are realized in buckypapersmaterials, which assists the production of high nanoscale fiber content(i.e., greater than 20 wt. %) buckypaper for high performance compositesmaterials.

In various embodiments, the films have an average thickness from about 5to about 100 microns thick with a basis weight (i.e., area density) ofabout 20 g/m² to about 50 g/m². In one embodiment, the buckypaper is athin film (approximately 20 μm) of nanotube networks.

Upon mechanical stretching of the buckypapers, the aligned nanoscalefiber films may have a tensile strength ranging from 20 MPa to 668 MPa.In other embodiments, the aligned nanoscale fiber films have a tensilestrength ranging from 200 MPa to 668 MPa. In yet other embodiments, thealigned nanoscale fiber films have a tensile strength greater than 668MPa.

In other embodiments, the aligned nanoscale fiber films have a Young'smodulus ranging from 2 GPa to 25 GPa. In still other embodiments, thealigned nanoscale fiber films have a Young's modulus ranging from 5 GPato 25 GPa. In yet other embodiments, the aligned nanoscale fiber filmshave a Young's modulus greater than 25 GPa.

Nanoscale Fiber Composites and Uses Thereof

Supporting Medium

Suitable supporting medium for use in stretching methods describedherein include PE, polypropylene, polymethyl methacrylate,polycarbonate, polyamide, or any other materials known in the art whichare compatible for use with nanoscale fibers. In certain embodiments,the supporting medium is capable of thermal decomposition so thatsubstantially all of or at least a portion of the supporting medium isremoved from the composite. In other embodiments, the supporting mediumis dissolved by a solvent.

Matrix Material

Composite materials are provided that comprise aligned nanoscale fibersand a supporting medium and/or a matrix material. In one embodiment, thecomposite materials may include a matrix material and the supportingmaterial may or may not be present in the composite. Suitable matrixmaterials include epoxy resins, phenolic resins, bismaleimide (BMI),polyimide, thermoplastic resins (e.g., nylon and polyetheretherketoneresins), and other polymers.

In certain embodiments, the matrix material may comprise a B-stage curedresin (e.g., an epoxy, a polyimide, a bismaleimide, a phenolic resin, acyanate, or a combination thereof) such that the composite materialcomprises a prepreg.

Composites

In certain embodiments, composites comprising mechanically stretchedbuckypaper have a nanoscale fiber concentration (loading) ranging from 5wt % to 62 wt %.

In certain embodiments, composites comprising mechanically stretchedbuckypaper have a tensile strength ranging from 620 MPa to 2,088 MPa. Inother embodiments, composites comprising mechanically stretchedbuckypaper have a tensile strength ranging from 620 MPa to 3,081 MPa. Inyet other embodiments, composites comprising mechanically stretchedbuckypaper have a tensile strength greater than 3,081 MPa.

In other embodiments, composites comprising mechanically stretchedbuckypaper have a Young's modulus ranging from 47 GPa to 169 GPa. Instill other embodiments, composites comprising mechanically stretchedbuckypaper have a Young's modulus ranging from 47 GPa to 350 GPa. In yetother embodiments, composites comprising mechanically stretchedbuckypaper have a Young's modulus greater than 350 GPa.

Embodiments of the aligned nanoscale fibers may also have tight packing,which is desirable for high loading in composites. Improvement of thealignment of the nanoscale fibers and loading is a desirable factortoward realizing the potential of nanotubes for high mechanical,electrical and thermally conductive applications in composites andelectronic devices. Thus, embodiments of the high-performance buckypapernanocomposites can be used for EMI shielding, thermal management andstructural materials applications. Immediate applications includecomposite applications for aircraft and thermal management forelectronic device package. Other applications may include lightningstrike protection, other lightweight structural materials applications,and electronic and energy applications, such as high-conducting thinfilm and powerful and efficient battery and fuel cell electrodes.High-performance buckypaper materials may also be used to developlightweight-conducting films and current-carrying materials forelectronic products.

The methods and compositions can be further understood with thefollowing non-limiting examples.

Example 1

BP/PE composite films were made using the following procedure: The BPwas sandwiched with PE film (two layers on both sides of the BP) in avacuum bag with 14.7 PSI vacuum pressure at 190° C. for 30 min. TheBP/PE composite was cooled in the vacuum bag to room temperature. Thebuckypaper used was slightly aligned (i.e., less than 20% degree ofalignment) MWNT film sheets from Nanocomp (Concord, N.H.). The aerialdensity of the buckypaper was 20-25 g/m². The PE film was clear plasticcling wrap that was manufactured by the GLAD® Products Company (CA94612), which had an area density of about 13.5±0.5 g/m².

The BP/PE composite films were stretched using the following steps: Thecomposite strip was mounted on the clamp of a Shimazu machine (AGS-JModel) and stretched by the machine at a speed of 1 mm/min to thedesired deformation or stretching ratio. An electric dryer (Dryer 00415,L&R® Manufacturing, NJ 07032) was used to heat the composite film toabout 60±10° C. during stretching. FIG. 3 shows the stretching setup andstretched BP/PE film used.

Before the stretched PE/BP composite films were “burned,” the stretchedBP composite film was treated in a vacuum bag at 90-100° C. for 10minutes to disorient the PE molecules and keep the nanotubes' alignment.Otherwise, the composite film would have shrunk at high temperaturesduring the burning process, thus affecting the quality of the stretchedBP. After being annealed, the BP/PE composite film was put into a tubefurnace (GSL-1600X, MTI Co. Richmond, Calif.) in a nitrogen or argonatmosphere at 550° C. for 1 hour to burn off the PE molecules in thecomposite film.

After burning off the PE material, the alignment of the nanotubes wasanalyzed through mechanical tests, Raman spectrum and electronicscanning microscope (SEM) experiments. FIG. 4 shows the degree ofalignment via Raman spectrum comparison of the stretched (30% stretchratio) and unstretched BP samples using a polarized Raman spectrometer(InVia, Renishaw, Inc). A high Raman G-band intensity ratio along thestretching direction (0 degree direction) indicated nanotube alignmentalong that direction. In particular, large angular changes of Ramanspectrum intensity (normalized) indicated more nanotube alignment in thestretched MWNT buckypaper.

Mechanical tests of neat buckypaper samples were also conducted toreveal the effects of stretching on mechanical properties. Samples withdifferent stretching ratios were tested using a Dynamic MechanicalAnalysis (DMA) machine (DMA 2980, TA Instrument Co.) under control forcemode to obtain tensile strength and modulus values. FIG. 5 shows typicalstress-strain curves, indicating noticeable property changes of thestretched BP samples. FIGS. 6 and 7 show the modulus and strengthincreased with the increase of the stretching ratios, respectively. Thebreaking strains decreased with the increase of the stretching ratios,indicating alignment improvement in the samples. The modulus of thepristine BP was about 4.67N/tex (GPa/(g/cc)), and the modulus of the BPsample with a stretching ratio of 40% went up to 16.96N/tex, achieving a3.6 fold increase. The results also showed the non-linear behaviors ofthe modulus and strength increase with an increase of the stretchingratio. A dramatic modulus increase was observed with stretching ratioshigher than 30%. For the strength increase, the same transition pointwas about 25% of the stretching ratio. However, there exists a peak ofstrength increase. A very high stretching ratio will lead to damage ofthe nanotube network, and the strength would decrease. The results showthat the strength of the neat buckypaper samples increased from0.143N/tex (pristine buckypaper) to the maximum of about 0.21N/tex atapproximately 30%-35% stretching ratio. The maximum strength increasewas only about 50% of the pristine sample performance, much less thanthe approximately 360% enhancement in the modulus.

FIG. 8 shows the SEM images of buckypapers with different stretchingratios. A significant improvement of nanotubes in the stretchedbuckypaper samples was observed, which further indicated theeffectiveness of the stretching process.

Stretched MWNT buckypaper/epoxy buckypaper composites were also producedand tested to determine the effectiveness of stretching on the compositeproperties. MWNT buckypaper samples with a 30% stretching ratio wereused after burning off the PE film. Epon 862 epoxy and curing agent W(diethyltoluenediamines) from E.V. Rubber Inc. were used as received forcomposite fabrication. The stretched BP samples were impregnated with anEPON 862 resin solution. The solution was 15 wt. %-20 wt. % epoxy, andthe curing agent W (weight ratio: 100:26.4) was in acetone. Afterevaporating the acetone in a vacuum oven at 70±10° C. for 30 min at avacuum degree lower than 1 Psi, from the resin impregnated buckypaper,ten layers of the impregnated buckypaper sheets were stacked togetherand hot pressed using a hot press machine Model 3925, from Carver Inc.The composite curing parameters were 177° C. for 2.5-3 hours at 20 MPapressure. Composite testing was conducted according to ASTM D 638-03standards, and the Shimazu machine was used. The typical stress-straincurves of the control (unstretched) and stretched composites are shownin FIG. 9, revealing that the tensile modulus and strength of thestretched 30% BP composites was 125.1 GPa and 1,056 MPa, respectively,which is much higher than the values of the control sample. Significantproperty improvements were achieved through the stretching process.

Example 2

A bismaleimide (BMI) resin (BMI resin 5250-4 from Cytec Inc.) was usedto make composites with improved thermal and mechanical properties. Inparticular, composites reinforced with BMI were fabricated with highcarbon nanotube loading (up to 60 wt %) BP. The BP used were randomlyoriented MWNT film sheets from Nanocomp (Concord, N.H.).

Mechanical Stretching Alignment

The random BP were stretched using a Shimadzu machine to align the CNTs.The crosshead speed during stretching was 0.5 mm/min in all stretchingexperiments. Since the randomly oriented CNT BPs had good stretchcharacteristics and strength, they could be stretched without using asupporting medium. The randomly oriented CNT BPs were stretched to astretch ratio up to about 40%. Attempts to stretch the BPs beyond a 40%stretch ratio led to breakage of CNT BP strips. The alignment degree wascharacterized by polarized Raman spectrum and small angle X-rayscattering.

Fabrication of BP/BMI Composites

The bulk BP/BMI composite samples were fabricated with a BP prepreg andhot press molding technique. First, the buckypaper was impregnated withBMI resin solution to make a prepreg with about 60±2 wt % CNTconcentration. The concentration of the BMI resin solution was adjustedto be less than 10 wt % to ensure a low viscosity for facilitatingimpregnation. The solvent used was acetone. Then six layers of BPprepreg were stacked together and cured with a 25 MPa hot-press processfollowing a curing cycle of 375° F. for 4 hours and 440° F. for 2 hours.The CNT weight percent of the final composite samples was controlled at60±1%. Using this method, high quality CNTs/BMI composites with highloading, good dispersion and alignment of CNTs in BMI matrix were made.

Carbon fiber polymer composites were also fabricated using these prepregand hot-press steps.

Characterization

Mechanical property tests of the BP/BMI composites were conducted by aShimadzu machine (AGS-J) according to ASTM D 638-03. The crosshead speedwas 1 mm/min with 20 mm gauge length. Fracture surfaces of the tensiletest specimens were coated with gold layer and observed using aelectronic scanning microscope (SEM: JEOL JSM-7401F USA, Inc.). The SEMimage after stretching showed that the CNTs are well aligned. DMA wasperformed on a DMA 800 machine (TA instrument Inc.) using film mode witha constant frequency of 1 Hz in a temperature from room temperature to400° C. with the heating rate of 5° C./min. The electrical conductivityproperties of the BP/BMI composites were characterized by a standardfour-probe method using a source meter.

Results and Discussions

Mechanical Properties and Fracture Surface Morphology

Four stretched, aligned CNT BP having different stretch ratios werecompared with randomly oriented CNT BP. The DMA static mode was used totest the mechanical properties of CNT BP before and after stretching.The stress-strain curves are shown in FIG. 10. The tensile stress of CNTBPs increased after stretching from 0.095 N/tex for randomly orientedCNT BP to 0.118 N/tex, 0.123 N/tex, 0.139 N/tex and 0.153 N/tex for 30%,35%, 38% and 40% stretched BP, respectively. Compared to tensile stress,the Young's modulus exhibited even more improvement afterstretching—from 0.346 N/tex for randomly oriented CNT BP to 4.150 N/tex,5.105 N/tex, 5.433 N/tex and 6.970 N/tex for 30%, 35%, 38% and 40%stretched BP, respectively.

Tensile stress-strain curves for CNTs/BMI composites are shown in FIG.11A, and the tensile strength and Young's modulus of these samples areshown in FIG. 11B. The effect of CNT orientation on their mechanicalproperties appeared to be significant. The BP/BMI composite showed amuch higher tensile strength and Young's modulus as compared toliterature reported values. The mechanical properties were noticeablyimproved with stretching ratio increase because of better CNT alignment.

The tensile strength and Young's modulus of 30% stretched BP/BMIcomposite samples were about 1639 MPa and about 124 GPa, respectively.When the BP was further stretched to 35%, the tensile strength andYoung's modulus of its composites was increased to about 1896 MPa andabout 146 GPa, respectively. The tensile strength and Young's modulus ofthe 40% stretched CNTs/BMI composites were 2088 MPa and 169 GPa,respectively.

FIG. 12 shows the comparison of mechanical properties between theCNTs/BMI composites with state-of-the-art unidirectional (UD) carbonfiber composites for aerospace structural applications. The mechanicalproperties of CNTs/BMI composites were normalized to 60 vol % with thedensity of CNTs 1.8 g/cm³. The randomly oriented, or pristine, BP/BMIcomposites were comparable to or lower than that of carbon fiber fabriccomposites for structural applications. After CNT alignment bystretching, the mechanical properties dramatically increased with thealignment degree of CNTs. When the CNT BP was stretched to 35%, theYoung's modulus exceeded the traditional carbon fiber composites. Whenthe CNT BP was stretched to 40%, the tensile strength is comparable tothat of unidirectional Toray T700 and IM7 carbon fiber composites andhigher than that of unidirectional Toray T300, AS4, AS7, and IM7 UDcarbon fiber composites.

The fracture surface morphology of the BP composites samples is shown inFIG. 13. It can be seen that only a few broken nanotubes (CNTs) wereobserved at the fracture surface of the 35% stretched BP/BMI composites,as shown in FIG. 13B. The CNTs formed large strips due to possibleself-assembly during fracture deformation, as shown in FIGS. 13B and13D. Such self-assembly may lead to lower stress concentration in CNTs,and hence improve composite properties. On the other hand, this alsoindicates that further improvement of alignment and interfacial bondingshould lead to more CNT breaks (i.e., more loading on the CNTs and thus,composite fracture due to CNT breaks rather than BMI fracture), whichcould lead to much higher mechanical performance.

Dynamic Mechanical Analysis

The DMA tests were conducted to confirm the Young's modulus and T_(g) ofBP/BMI composites.

FIG. 14A shows the comparison of storage modulus of the CNT_(S)/BMIcomposites. The storage modulus of the composites was 55.4 GPa, 123.3GPa, 146.7 GPa and 171.7 GPa, respectively corresponding to the CNTs/BMIcomposite, 30% stretched CNTs/BMI composite, 35% stretched CNTs/BMIcomposite and 40% stretched CNTs/BMI composite. The storage modulus ofthe CNTs/BMI composites was consistent with Young's modulus tested bytensile testing. The T_(g) determined by Tan Delta is 269.98° C.,266.77° C., 259.76° C. and 256.70° C., respectively, corresponding tothe CNTs/BMI composite, 30% stretched CNTs/BMI composite, 35% stretchedCNTs/BMI composite and 40% stretched CNTs/BMI composite, seen FIG. 14B.Although the CNT loading was as high as approximately 60 wt %, the T_(g)was fairly consistent, which showed that the introduction of highloading CNTs does not reduce the cross-linkage of BMI resin as large asobserved in epoxy resin systems. Such high T_(g) and good thermalstability of the BP/BMI composites are desirable for high temperaturestructural applications. The reduction of the Tan Delta areas of thestretched MWNT composites could imply more molecular interactionsbetween MWNTs and BMI due to more spreading of the MWNT ropes and largeinterface areas.

Electrical Properties

The electrical properties comparison of the CNT/BMI composite is shownin FIG. 15. Due to the high loading of CNTs in the composites, theelectrical conductivity was high. For the randomly CNT/BMI composite,the electrical conductivity reached 915 S/cm, which was higher than thatof reported results of carbon fiber reinforced composites. The improvedCNT loading and the contact of CNTs may have been the two reasons forthis increase. For the stretched CNT/BMI composites, the electricalconductivity along the CNTs direction (σ_(//)) was much higher: 1774S/cm, 2954 S/cm and 5473 S/cm corresponding to 30% stretched CNTs/BMIcomposite, 35% stretched CNTs/BMI composite and 40% stretched CNTs/BMIcomposite, respectively. The reason for this increase was likelywell-oriented CNTs in the BMI matrix seen FIGS. 13C and 13D. Such highcomposite conductivity enables EMI shielding and lightning strikemultifunctional applications of structural composites.

Conclusions

High mechanical performance nanotube buckypaper/BMI composites withcomparable mechanical properties to the state-of-the-art highperformance carbon fiber composites were successfully demonstrated. Thiswas a desirable milestone for developing high performance nanotubecomposites for structural applications.

Example 3

Slightly aligned (i.e., less than 20% degree of alignment) MWNT filmsheets from Nanocomp Technologies Inc. were mechanically stretched usinga Shimadzu machine (AGS-J, Shimadzu Scientific Inc., Japan) to enhancenanotube alignment. The randomly dispersed MWNT sheets manufactured byNanocomp Technologies Inc. (Concord, N.H.) include millimeter-long andsmall-diameter (˜3-8 nanometers) MWNTs with a range of 2-5 walls,providing an aspect ratio up to 100,000.

The crosshead speed during stretching was 0.5 mm/min. There was noobserved retraction after stretching. The stretching ratio of the MWNTsheets was calculated using Equation 1. MWNT sheets were processed atthree stretch ratios (30%, 35%, and 40%) for composite fabrication. Thecomposite samples had approximately 60 wt % nanotube weight fraction orloading. For 40%-stretched CNT sheet (i.e., the post-stretch sheet was40% longer than the pre-stretch sheet), the degree of alignment of theCNT sheet was dramatically improved.

To further understand and quantify the effects of the stretch ratio onthe alignment degree, polarized Raman scattering tests were conductedand the alignment degree was calculated. Polarized Raman intensity ofthe G-band was measured as a function of angle between laser polarizeddirection and nanotube alignment direction or stretch axis. G-band Ramanintensity showed the maximum if the polarization was parallel to thestretched axis (θ=0) and was at a minimum at the perpendicular direction(θ=90).

Theoretical Raman intensity change simply follows cos⁴ θ versus nanotubeorientation angle (θ), A two-dimensional distribution function was usedto describe nanotube orientation distribution in the MWNT sheets. Then,the best fitting curve of the Raman intensity versus orientation anglecan be obtained, as shown in FIG. 16. From the trend of the best fittingcurve, the near perfect alignment (more than 95% nanotubes aligned alongstretch direction) at an approximate 50% stretch ratio was predicted.However, the actual highest stretch ratio could only reach about 40%,Attempts to stretch CNT sheets over 40% were not successful due tonanotube network breakage beyond what the nanotube aspect ratio couldhandle.

Both randomly dispersed (i.e., as-received) and stretched MWNT sheetswere used to make MWNT/BMI resin matrix composites. Aerospace-grade BMIresin 5250-4 (Cytec. Inc) was used as the matrix resin. First, the MWNTsheets were impregnated with a BMI resin solution of the samecomposition as used in Example 2 to make individual MWNT prepreg sheetswith approximately 60±2 wt % nanotube concentration (i.e., loading). Theprepregging process involved solution impregnation under 2-5 MPapressure. The concentration of BMI resin in the solution was adjusted toensure low viscosity for facilitating impregnation. The residual solvent(i.e., acetone) was removed from the MWNT/BMI prepreg using a vacuumoven at 80° C. for 2 hours, resulting in the BMI/MWNT prepreg.

Second, six layers of the MINT sheet prepregs were stacked together andcured by the hot-press process under 5-20 MPa pressure following thecuring cycle of: 375° F. (190.5° C.) for 4 hours and then 440″F(226.7″C) for 2 hours. The weight percentage of the nanotubes wasdetermined by the weight of total amount of the MWNT sheets used duringcomposite fabrication. The MWNT weight ratio in the final composite wascontrolled to the range of 60±2 wt %. The weight fractions of therandomly oriented, 30%-, 35%-, and 40%-stretched MWNT composite sampleswere 60 w %, 61.7 wt %, 60.5 wt % and 61.7 wt %, respectively. Thedensities of MWNT and BMI resin were 1.8 g/cm³ and 1.25 g/cm³,respectively. The measured density values of the composite samples are1.525, 1.536, 1.530 and 1.536 g/cm³ respectively.

Hence, the calculated void volume fractions were 0.367 vol %, 0.286 vol%, 0.227 vol %, 0.286 vol % for the samples, respectively. Forcomparison purposes, the void volume fraction requirement ofconventional structural carbon fiber composites was less than 2 vol %.Hence, the samples had good \vetting between MWNTs and BMI resin,indicated by the low void volume content.

Characterization Mechanical properties tests were conducted using aShimadzu machine with crosshead speed of 1 mm/min and gauge length of 20mm under room temperature. The strain was recorded by Shimadzunon-contact video extensometer DVE-201. The specimens were cut into adog-bone shape with a length of 35 mm and a working length of 20 mm andthickness of 60 mm according to ASTM D638. After the tensile tests, thefracture surface morphology of the specimens was coated with a goldlayer and observed using an electronic scanning microscope (JEOLJSM-7401F USA, Inc.). DMA was performed on a DMA 800 machine (TAinstrument Inc.) using the film mode with a constant frequency of 1 Hzfrom room temperature to 400° C. with a heating rate of 5° C./min.

Load carrying along the alignment direction was seen in thepost-stretching samples. The mechanical properties of the neat (i.e.,without polymer binders or resin) MWNT sheets of different stretchratios were measured, as shown in FIG. 17. The tensile strength at breakand Young's modulus of a randomly dispersed CNT sheet (the controlsample) were approximately 205 MPa and 1.10 GPa, respectively. Duringstretching, the MWNTs self-assembled and aligned themselves along theload direction.

Along the alignment direction, the mechanical properties were alsoimproved. The tensile strengths increased to 390 MPa, 508 MPa, and 668MPa for the 30%, 35%, and 40% stretched samples, corresponding to 90%,148% and 226% improvements, respectively. The post-stretch Young'smodulus measurements along the alignment direction showed even moredramatic improvements, from 1.10 GPa for the randomly dispersed sheet(pre-stretch) to 11.93 GPa, 18.21 GPa, and 25.45 GPa, respectively,showing of 10-, 16-, and 22-fold improvements. Compared to other CNTsheets, the MWNT sheets used in this study resulted in moreentanglements that maintained the integrity of the nanotube networks dueto a large aspect ratio. As a result, their mechanical properties andcreep resistance were relatively high.

FIG. 18 shows the electrical conductivity measurements of the neat MWNTsheets paralleled to the alignment direction. The electricalconductivity was higher with the increased stretch ratio. Electricalconductivity (σ_(//)) paralleled to the alignment direction increasedfrom 420 S/cm in the pre-stretched CNT sheets to 600 S/cm in the CNTsheet with 40% stretch ratio. The electrical conductivity of thestretched sheets was not very high likely because the neat MWNT sheetswere still porous.

Mechanical Properties and Fracture Morphology of MWNT Sheet/BMINanocomposites

The MWNT/BMI composites demonstrated a relatively large 2.0-2.5% failurestrain as compared to carbon-fiber-reinforced composites, which aretypically in the range of 0.6-1.8%. The large tensile strains exceeded2.0% because of the CNTs' intrinsic flexibility, high failureelongation, and high deformability of the MWNT networks in the sheets,as shown in FIG. 17. The tensile strength of the randomly dispersedMWNT/BMI composite (the control sample) was approximately 620 MPa, andthe Young's modulus was 47 GPa, After stretching to improve alignmentand nanotube packing, the mechanical properties dramatically increased.The tensile strength and Young's modulus of the 30%-stretched CNT/BMIcomposite were 1,600 MPa and 122 GPa, respectively. When the stretchratio increased to 35%, the tensile strength and Young's modulusincreased respectively to 1,800 MPa and 150 GPa. The tensile strengthand Young's modulus of the 40%-stretched MWNT/BMI composite were as highas 2,088 MPa and 169 GPa, respectively.

The total number of nanotubes in the axial tensile directiondramatically increased (shown in FIG. 16) with an increase in degree ofalignment. Efficiency of both load carrying and transfer for the alignednanotubes in the axial tensile direction was significantly enhanced,leading to dramatically higher mechanical properties. For example, the40%-stretched sample had an alignment degree of the MWNTs along theaxial direction of ˜0.8 (seen in FIG. 16), which meant about 80% of thenanotubes were probably aligned along the stress direction to carry aload when the tensile stress was applied. Interfacial bonding betweenthe CNTs and resin matrix was also a relevant factor. Typically, the CNTand polymer chain are weak due to CNTs' atomically smooth surfaces.

FIG. 19A shows the fracture surface morphology of a 40%-stretchedspecimen after tensile tests. The MWNTs were peeled off as very thin andtransparent films. BMI resin was suspected to coat on the nanotubebundle surface because high nanotube concentration and no bulk neatresin fractures were observed. This peeled off failure mode was not aresult of individual MWNT sheets sliding to each other because thethickness of individual sheet is more than 10 μm. The thickness of thepeeled off thin films are much less than 100 nm (FIG. 19C). Many stretchdeformations of the MWNT/BMI thin films were observed, indicatingeffective load transfer between the MWNTs and BMI resin matrix in thecomposites. The evidence of MWNT slippage failure mode also was seen inFIGS. 19A and 19B. The MWNTs were pulled out from the composites andbecame very stretched strips with obvious diameter change and sharpbreaks at the end due to MWNT slippage within the bundles. Furthermore,although the resultant composites showed record-high mechanicalproperties, almost no broken nanotubes were seen—evidence that the fullpotential of CNTs' strength has yet to be completely realized. Theformed CNT/BMI thin films were transparent (seen in FIG. 19B),indicating that the thickness consisted of only a few layers ofwell-spread nanotubes. The spreading of nanotube bundles meant the MWNTropes morphed from original round and large-diameter shapes into flatthin-film shape during to mechanical stretching and prepreggingprocesses. Such thin-film strips of nanotube assemblies have morenanotubes at the outmost layer to interact with other nanotubeassemblies and resin matrix to achieve good load transfer. The MWNTswere also well-aligned along the loading direction, which helped inrealizing good load carrying. Such unique microstructures of theMWNT/BMI nanocomposites were results of the mechanical stretching andprepregging under pressure.

Thermal Mechanical Performance of MWNT Sheet/BMI Nanocomposites

DMA tests were conducted to confirm measured Young's modulus and measureglass transition temperature (T_(g)) values of the MWNT/BMI samples. Theresults were similar to those see in the composites of Example 2. Thestorage modulus measurements of the composites were 55 GPa, 123 GPa, 147GPa, and 172 GPa for the pre-stretched sample (the control sample),30%-, 35%-, and 40%-stretched MWNT/BMI samples, respectively. Thestorage moduli of the CNT/BMI composites were consistent with theYoung's modulus values in the tensile testing.

Electrical Conductivity of MWNT/BMI Nanocomposites

The electrical conductivities of the MWNT/BMI composites were alsosimilar to those of the composites made in Example 2.

Conclusions

In summary, high mechanical and electrical properties of MWNT/BMIcomposites were realized. The coupling effects of MWNTs an averagelength of at least 1 millimeter, mechanical stretching, and prepreggingunder high pressures led to high loading, good alignment, and enhancedload transfer. These factors led to mechanical property improvements.Additionally, successful dispersion of the nanotube ropes intospread-out extra-thin films led to better contacts among MWNTs, givingrise to effective load transfer and enhanced electrical conductivity.Integration of the high mechanical properties and electrical conductanceindicated these MWNT/BMI composites will lead to excellent materials forlightweight composite conductors for a wide range ofmultifunctional/structural applications.

Example 4

Functionalized CNT sheets were used to reinforce BMI composites. Themechanical properties of the resultant CNT sheet/BMI composites werenormalized to 60 vol % nanotube volume content and compared with theunidirectional (UD) carbon fiber composites. These compositesdemonstrated mechanical properties beyond aerospace—grade unidirectionalcarbon fiber composites for structural applications.

Materials and Functionalized CNT Sheet/BMI Nanocomposite Fabrication

Randomly oriented CNT sheets (supplied by Nanocomp Technologies Inc.)were mechanically stretched using an AGS-J Shimadzu machine tosubstantially improve nanotube alignment as described in Example 3. Theresin system used was Cytec's BMI 5250-4 resin, which contains threecomponents, 4, 4′-bismaleimidodiphenylmethane, o,o′-diallyl bisphenol Aand BMI-1, 2-tolyl. According to a phenol-epoxy curing mechanism, theactive epoxy groups can react with hydroxyl groups of o,o′-diallylbisphenol A. Hence, epoxidation functionalized CNTs were used to realizecovalent bonding with BMI resin matrices. This functionalization methodwas suitable for tailoring the degree of functionalization (DOF) using agentle reaction condition to avoid damage of preformed nanotubealignment and sheet structural integrity.

Peroxide acid (m-chloroperoxybenzoic acid, m-CPBA) was used to treatsingle-walled carbon nanotubes (SWNT) and introduce an epoxy ring on thestructure of the SWNTs. Both randomly dispersed and aligned CNT sheetswere treated with a m-CPBA solution prepared in the same manner as them-CPBA solution of Example 3 to realize a tailored 4% functionalizationdegree to minimize CNT structure damage and composite mechanicalproperty degradation. Specifically, the aligned CNT sheets were placedin a m-chloroperoxybenzoic acid (m-CPBA)/dichloromethane solution forepoxidation functionalization, and then washed using dichloromethane toremove residual m-CPBA. The functionalized CNT sheets were placed intothe vacuum oven at 80° C. for 30 min to evaporate the residualdichloromethane.

Then, CNT sheets were impregnated with BMI 5250-4 resin solution to makeindividual CNT prepreg sheets with approximately 60 wt % nanotubeconcentration or loading. The BMI resin solution was prepared in thesame manner as the BMI resin solution described in Example 3. Theprepregging process was a solution impregnation process. The residualsolvent (acetone) was removed under 80° C. in the vacuum oven for 2hours to make BMI/CNT sheets prepreg. Six prepreg layers were stackedtogether and cured by the hot-press with 5-20 MPa pressure following thecuring cycle: 375° F. for 4 hours and then 440° F. for 2 hours. The CNTweight fraction in the final composites was 60±2 wt %.

Characterization:

Mechanical properties test were conducted using a Shimadzu machine withcrosshead speed of 1 mm/min and the gauge length of 20 mm under roomtemperature. The strain ratio was recorded by Shimadzu non-contact videoextensometer DVE-201. The specimens were cut into dog-bone shape with alength of 35 mm and thickness of 60 μm according to ASTM D638. Thetypical tensile stress-strain curves of functionalized CNT sheet/BMIcomposite are shown in FIGS. 20-22. After the tensile tests, thefracture surface morphology of the specimens was coated with a goldlayer and observed using an electronic scanning microscope (JEOLJSM-7401F USA, Inc.). DMA was performed on a DMA 800 machine (TAinstrument Inc.) using the film mode with a constant frequency of 1 Hzfrom room temperature to 400° C. with a heating rate of 5° C./min. Theelectrical conductivity of the functionalized CNT sheet/BMI compositeswas measured using the four-probe method.

As shown in FIG. 23, the mechanical properties of pristine 40% stretch(stretched to a 40% strain to increase nanotube alignment) CNT sheet/BMIcomposites achieved the mechanical properties of standard UD carbonfiber reinforced composites, such as AS4 and T300 carbon fibercomposites. After functionalization, the mechanical properties offunctionalized 40% stretch alignment CNT sheet/BMI composites wereimproved to an even higher level. The Young's modulus exceeded that ofhighest-modulus carbon fiber composites, such as M60J epoxy composite,and the tensile strength was 15-20% higher than that of high-strengthT1000G epoxy composites.

FIG. 24 is a graph of ATR-FTIR spectra of pristine CNTs, functionalizedCNTs, and pristine and functionalized aligned (40% stretch) CNTsheet/BMI composites. The peak at 1210 cm⁻¹ was attributed to epoxy ringgroups, which confirms the epoxy group successfully attached to the CNTstructure. After curing with BMI resin (see Trace d), the peak at 1210cm⁻¹ dispeared, which indicated the epoxy ring group reacted with BMIresin. The FTIR spectra of pristine CNT sheet/BMI composite is shown inTrace c. Both FTIR spectra were almost same, which further confirms theepoxy rings on the CNT structures reacted to form covallent bonding withthe BMI resin matrix.

FIG. 25 shows the proposed reaction mechanism. The epoxide groups offunctionalized CNT first reacted with o,o′-diallyl bisphenol A inaccordance with the mechanism of epoxy-phenol reaction. Then, thederivative reacted with the other two BMI components to form the threedimensional crosslinked structures through ENE and Diels-Alderreactions. The formation of carbon-oxygen bonds between CNT and BMIresin dramatically enhanced the interfacial bonding, and hence the loadtransfer efficiency was improved after functionalization.

The curing mechanism was also studied using a Raman spectrometer. Theintensity ratio of disorder band (D band at ˜1310 cm⁻¹) with G band(˜1580 cm⁻¹) of the functionalized CNT increased, which indicates theformation of epoxy rings on the structure of the CNTs, as shown in FIG.26. The R-value (I_(D)/I_(G)) of pristine CNTs was 0.13. Afterfunctionalization, the I_(D)/I_(G) value increased up to 0.41. In thepristine CNT sheet/BMI composite, the I_(D)/I_(G) value increased to0.23 due to the coupling effect of CNTs and BMI crosslinked structure.For functionalized CNT sheet/BMI composite, the I_(D)/I_(G) furtherincreased up to 0.62, which further indicates stronger interactions,possibly due to the formation of chemical bonding between thefunctionalized CNT with BMI resin.

FIG. 27 shows the typical stress-strain curves of CNT sheets reinforcedBMI nanocomposites along the nanotube alignment direction. For pristinerandom CNT sheet reinforced BMI nanocomposites, the tensile strength andYoung's modulus dramatically increased as the alignment degreeincreased. The degree of nanotube alignment had a significant impact onthe mechanical properties. The results show the degree of CNT alignmentcan reach as high as 80% along the stretching or alignment directionwhen the CNT sheets were stretched to a 40% strain. The tensile strengthand Young's modulus of the resultant CNT sheet/BMI composites were ashigh as 2,088 MPa and 169 GPa, respectively.

After functionalization to introduce epoxy groups on the CNTs and thencovalently bonding with the BMI resin matrix, the mechanical propertiesof the resultant nanocomposites were dramatically improved. The tensilestrength and Young's modulus of functionalized random CNT sheet/BMInanocomposites reached up to 1,437 MPa and 124 GPa, respectively, whichis two times greater than that of pristine random CNT sheet/BMInanocomposites previously reported. For functionalized 30% stretchalignment CNT sheet/BMI nanocomposites, the tensile strength and Young'smodulus reached up to 2,843 MPa and 198 GPa, which is a 78% and 62%improvement above that of the pristine 30% stretch alignment CNTsheet/BMI nanocomposites. For functionalized 40% stretch alignment CNTsheet/BMI nanocomposites, the tensile strength and Young's modulusreached up to 3,081 MPa and 350 GPa, which are 48% and 107% improvementsover that of pristine 40% stretch CNT sheet/BMI nanocomposites. However,the failure strain of functionalized CNT sheet/BMI nanocompositesdecreased sharply, as shown in FIG. 27A. The failure strain offunctionalized 40% stretch alignment CNT sheet/BMI nanocompositesdropped to 0.95%. This may be due to two possible reasons: (1) theformation of covalent bonding significantly reduced nanotube pullout andrestricted nanotube network deformation capability and (2) possiblenanotube structural damage due to the functionalization resulted in aloss of certain degree of ductility of the CNTs. Therefore, degree offunctionalization should be examined and optimized to improve strengthand modulus without sacrificing failure strain. Here, the degree offunctionalization was adjusted to 4% to minimize CNT damage and failurestrain reduction of the composites.

FIGS. 29A-B show the fracture surface morphology of a functionalized 40%stretch alignment composite after tensile testing. Rather than peelingoff as seen in the pristine CNT sheet/BMI samples previously reported,it can be seen that BMI resin and aligned CNT layers adhered well due togood interfacial bonding. Although the interfacial bonding and loadtransfer efficiency were dramatically improved with this chemicalfunctionalization, resulting in the high mechanical properties exceedingthat of the state-of-the-art aerospace-grade unidirectional carbon fibercomposites, many CNT slippage and pulled-out modes were still observed.Also, most of nanotubes were not broken after tensile testing, whichimplies the full potential of CNTs' strength has yet been achieved.

FIGS. 30A-B show dynamic mechanical analysis (DMA) results. Table 1shows the storage modulus of the samples.

TABLE 1 Storage modulus T_(g) Specimen (GPa) (° C.) Pristine random CNTsheet/BMI composite 55 269.98 Functionalized random CNT sheet/BMIcomposite 122 262.67 Pristine 30% stretch CNT sheet/BMI composite 123266.77 Functionalized 30% stretch CNT sheet/BMI 203 241.80 compositePristine 40% stretch CNT sheet/BMI composite 172 256.70 Functionalized40% stretch CNT sheet/BMI 354 247.44 compositeThe T_(g)s of all CNT sheet/BMI composites dropped due to theintroduction of high loading of CNTs, which possibly reduced thecrosslink density of the BMI resin matrix. Compared with pristine CNTsheet/BMI composites, the T_(g) of functionalized CNT composites furtherdropped, which may be due to the epoxy groups of functionalized CNTsreacting and consuming some functional groups of BMI resin, and hencefurther reducing crosslink density. However, the T_(g) drop of thefunctionalized CNT/BMI composites was only 23° C., and the compositesstill had a T_(g) of 247° C. for high temperature applications. Anotherside effect of chemical functionalization of CNTs is degradation ofelectrical conductivity. Usually, chemical functionalization will damageoriginal CNT electronic structure and lower the electrical conductivity.In this Example, the degree of functionalization was at a lower level,4%, to limit electrical conductivity degradation. The electricalconductivities of the functionalized CNT composites only show a smallreduction, less than 5%, due to the lower degree of functionalization.Conclusions

An epoxy group was introduced on CNT structures through epoxidationfunctionalization, and demonstrated high performance for the CNTsheet/BMI composites, which was beyond the state-of-the-art highstrength and high modulus unidirectional carbon fiber composites forstructural applications. The limited effect of CNT functionalization onT_(g) and electrical conductivity was observed due to a tailored lowdegree of functionalization. The results demonstrate great potential forutilizing CNTs to develop the next generation high-performancecomposites for wide structural and multifunctional applications.

Example 5

Development of high mechanical properties of CNT reinforced epoxycomposites was achieved by tailoring the DOF and improving alignment ofCNTs an average length of at least 1 millimeter. The resultantcomposites showed an unprecedented integration of high strength andmodulus, and large failure strain, compared to the state-of-the-artcarbon fiber reinforced composites.

Randomly oriented CNT sheets supplied by Nanocomp Technologies Inc. weremechanically stretched using an AGS-J Shimadzu machine to substantiallyimprove nanotube alignment. The aligned CNT sheets were placed inm-chloroperoxybenzoic acid (m-CPBA)/dichloromethane solutions forepoxidization functionalization, and then washed using dichloromethaneto remove residual m-CPBA. The functionalized CNT sheets were placedinto a vacuum oven set at 80° C. for 30 min to evaporate the residualdichloromethane. Finally, the CNT sheets were impregnated with a 10 wt %epoxy resin solution in acetone to make individual CNT prepreg sheetswith approximately 60% nanotube concentration or loading by weight. Theconcentration of epoxy resin in the solution must be adjusted to ensurelow viscosity for facilitating impregnation. Six prepreg layers werestacked together and cured by the hot-press with approximately 1 MPapressure following the curing cycle: 200° F. for 30 min and then 350° F.for 4 hours. The CNT weight concentration or loading in the finalcomposite samples was controlled in the range of 60±2 wt %.

Millimeter-long (1-2 millimeter) nanotubes used in this example were inthin sheets (20-25 μm), provided by Nanocomp Technologies. Epoxidegroups were introduced on the structure of CNT to directly functionalizethe CNT sheet materials through epoxidation functionalization, as shownin FIG. 30A. Epoxide groups created on the CNTs were very active andparticipated in the curing reaction of epoxy resin to realize covalentlybonding between the CNTs and epoxy resin matrix. The proposed reactionmechanism between the functionalized CNTs and epoxy resin matrix isshown in FIG. 30A. The epoxy ring group was first introduced throughfunctionalizing CNT sheets in m-CPBA/CH2Cl2 solutions. Then, the epoxyring groups on the CNTs reacted with curing agent-diethyltoluenediamine(DETDA). Finally, derivatives reacted with the Epon 862 molecules toform the three dimensional crosslinked structures through theDiels-Alder reaction.

DOF of the functionalized CNTs is defined as the ratio of the number ofcarbon atoms directly connected with oxygen atoms to the total number ofcarbon atoms of the CNT. To tailor the DOF values,m-CPBA/dichloromethane solutions of 0.5%, 1%, 2%, 5% and 10% by weightconcentrations were made. The functionalization was conducted at roomtemperature (22-25° C.) by varying reaction times from 10 minutes to 30hours. The CNT sheets were immersed into the solution for variousperiods of times, and removed to complete the functionalization withoutdamaging sheet structural integrity. The DOF values were determined bythe thermogravimetric analysis (TGA) in the range of 50-800° C. undernitrogen atmosphere.

Mechanical properties test were conducted using a Shimadzu machine witha crosshead speed of 1 mm/min and the gauge length of 20 mm under roomtemperature. The strain ratio was recorded by Shimadzu non-contact videoextensometer DVE-201. The specimens were cut into dog-bone shapes atlengths of 35 mm and 60 μm thick, in accordance with ASTM D638. Afterthe tensile tests, the fracture surface morphology of the specimens wascoated with a gold layer and observed using an electronic scanningmicroscope (JEOL JSM-7401F USA, Inc.). The pristine aligned CNT sheetreinforced epoxy composite was cut perpendicular to the CNT alignmentdirection using Leica EM UC6/FC6 ultramicrotome (German) and observed byhigh resolution transmission electron microscopy Tecnai F30 (Philips,Holland).

Results and Discussion

FIG. 31 shows the curves of DOF versus functionalization time and m-CPBAconcentrations. For all cases, the DOF values initially increasedrapidly with the reaction time and then reached an almost constantvalue. With the same treatment time, the DOF increased with the increaseof m-CPBA concentration, indicating the desired DOF can be accuratelytailored through adjusting reaction time and m-CPBA solutionconcentration. The goal of introducing the epoxy rings on the structuresof CNTs is to facilitate creating covalent bonding betweenfunctionalized CNTs and epoxy resin matrix.

FIG. 32A shows the attenuated total reflection Fourier transforminfrared (ATR-FTIR) spectrum comparison to verify the formation andreaction of the epoxide groups on functionalized CNTs ((a) pristine CNT,(b) functionalized CNT, (c) functionalized CNT sheet/epoxy composites,(d) pristine CNT sheet/epoxy nanocomposites and (e) cured neat epoxyresin.). Compared with pristine CNTs, the peak of 1210 cm-1 offunctionalized CNTs was assigned to the carbon oxygen stretchingfrequency of epoxide moiety as seen in Trace b. After curing with epoxyresin, this peak became smaller, showing that the epoxy ring groups onthe CNT reacted with epoxy resin, as seen in Trace c. The ATR-FTIRspectra of pristine CNT sheet/epoxy composite and pure epoxy resin areshown as Traces d and e. The peak of 1210 cm⁻¹ still existed in thepristine CNT sheet/epoxy composites due to residual epoxy group of Epon862 (epoxy resin matrix), same as Trace e of pure cured epoxy resin withthe same curing cycle.

Raman spectrometer was used to verify the proposed reaction mechanism.As shown in FIG. 32B, the R-value (ID/IG) of pristine CNT of 0.13indicates the quality of CNT is very good with a lower defect density((a) pristine CNT, (b) functionalized CNT, (c) pristine CNT sheet/epoxycomposite and (d) functionalized CNT sheet/epoxy composite). Afterfunctionalization, the ID/IG value increased up to 0.41, which indicatesepoxy rings formed on the structures of the CNT. For the pristine CNTsheet/epoxy composite, the ID/IG value increased to 0.30 due to thecoupling with cured epoxy crosslinked networks. For functionalized CNTsheet/epoxy composites, the ID/IG value further increased up to 0.99,which further indicates much stronger interactions between the CNTs andresin matrix due to the formation of chemical bond between thefunctionalized CNT with the epoxy resin matrix.

To further confirm the proposed reaction mechanism, the high resolutiontransmission electron microscopy (HRTEM) was conducted to observe thenanotube surface structure before and after functionalization, shown inFIG. 33A (pristine double-walled nanotube) and FIG. 33B (functionalizeddouble-walled nanotube). Most of nanotubes used in this example aredouble-wall nanotubes. After functionalization, the epoxide groups areattached on the outside wall which results in the roughness of nanotube,seen FIG. 33B.

To study the effect of different DOF on the mechanical properties ofnanocomposites, the DOF values of random CNT sheets were tailored to 4%,10% and 18%. FIG. 34 shows the mechanical properties of resultantnanocomposites. For the pristine random CNT sheet nanocomposites, thetensile strength and Young's modulus were 851 MPa and 45 GPa,respectively. After functionalization, the Young's modulus of CNT sheetnanocomposite increased. However, for all three different DOFs, theYoung's modulus was almost the same at 80 GPa. The effects ofinterfacial bonding enhancement between nanotubes and epoxy resin onload transfer efficiency may be at the same level for all three cases.However, the tensile strength of resultant nanocomposites with higherDOF values decreased, which indicates the high DOF damages the CNTstructure and degrades the CNT mechanical properties. The 4% DOF islikely adequate to substantially enhance load transfer between epoxyresin and functionalized CNTs without large strength degradation in theresultant nanocomposites.

To quantify load transfer efficiency improvement, a DOF-load transferefficiency model was proposed. The modified rule of mixtures (ROM)equation is used for predicting properties of discontinuous short fiberreinforced polymer composite, which assumes a perfect load transferefficiency between fibers and resin matrix. That is not true for CNTreinforced nanocomposites, as evidenced by many CNT pullout withoutbreaks and very low mechanical performance. Thus, the modified the ruleof mixture was used to consider load transfer efficiency effect, asshown in Equation (2).E _(c)=η₀·η_(L)·η_(B) ·V _(f) ·E _(f)+(1−V _(f))·E _(m)  Equation 2.where E_(c), E_(m), and E_(f) are Young's moduli of the resultantcomposites, matrix and fiber, respectively. V_(f) is the volume fractionof the CNTs. The orientation factor, η₀, was introduced to account forfiber orientation effect, which equals to 1 for fully aligned fibers.For randomly oriented fibers, the η₀ value was 0.33. The lengthefficiency factor, η_(L), was introduced to account for the efficiencyof load transfer from the matrix to the fibers due to aspect ratioeffect. η_(L) can vary between 0 and 1.

In this example, the length of the CNTs was approximately at themillimeter level, which is much larger than the diameters (3-8 nm) ofthe CNTs; therefore, η_(L) as 1. Herein, the interfacial loadingtransfer efficiency factor, η_(B), was defined and used to account forload transfer efficiency determined by interfacial bonding qualitybetween fiber and matrix. Equation (2) was changed into a logarithmicform to obtain Equation (3)lg<(E _(c)−(1−V _(f))=lg(η_(B))+lg(η₀·η_(L))÷lg(V _(f) ·E_(f))  Equation 3.Assuming η₃ is a function of DOF, then utilizing the results shown inFIG. 34, the curve of lg(E_(c)−(1−V_(f))·E_(m)) versus DOF, as shown inFIG. 35A. Through logistic fitting, the relationship between η_(B) andDOF directly can be shown, as seen in Equation (4) and FIG. 35B.

$\begin{matrix}{\eta_{B} = {10^{\frac{- 0.2182}{1 + {6.87 \times 10^{5} \times {({DOF})}^{3.3}}}}.}} & {{Equation}\mspace{14mu} 4}\end{matrix}$If DOF=0 and the η_(B,0)=0.605, which means the load transfer efficiencyinduced by nonbinding interfacial interactions is only 60.5% for thepristine CNT sheet of millimeter long nanotubes. If DOF=0.04, thenη_(B,0.04)=0.972, which means the load transfer efficiency is adequate.It also shows that η_(B) dramatically increased with the increase of DOFvalues at the beginning, then tended to become constant and saturated,which was in good agreement with other simulation results. Nanotubealignment is another factor to consider in realizing high mechanicalproperties as previously discussed. The sheets of randomly oriented longCNTs were stretched to about 40% strain to realize an alignment of ˜80%of the CNTs along the stretch direction, as determined by polarizedRaman analysis. The cross-section of random and aligned CNT sheets areshown in FIG. 36A (random) and FIG. 36B (aligned). After stretching,most nanotubes assembled along the stretching direction very well whichfurther verified ˜80% alignment degree determined by polarized Ramananalysis.

The highly aligned CNT sheet was further functionalized with a tailoredDOF of 4% to achieve a better performance of CNT reinforced epoxycomposites. FIG. 37 shows the typical stress-strain curves of CNT sheetreinforced epoxy nanocomposites with/without alignment andfunctionalization. After functionalization, the tensile strength andYoung's modulus of the random CNT sheets nanocomposites increased to1333 MPa and 80 GPa, respectively, as shown in FIG. 38. Such performanceis comparable to carbon fiber fabric composites. It is worth noting thatthe tensile failure strain of the pristine random CNT sheetnanocomposites reached 8.21%, which is much higher than that (3.5-5%) ofconventional carbon fiber fabric composites. Two possible reasons areattributed to this: (1) the pure randomly oriented CNT sheets have gooddeformation ability due to entanglements and slippages in the randomlyoriented networks of long CNTs; and (2) possible interface slippagebetween CNT and resin matrix can allow large deformations of the CNTnetworks within the composites. After functionalization, the interfacialbonding were dramatically enhanced due to the formation of chemicalbonding between CNTs and epoxy resin, which greatly constrains theslippage between CNT and epoxy resin and result in the low failurestrain of resultant nanocomposites.

The tensile strength, Young's modulus and failure strain of the alignedCNT composites were 2,375 MPa, 153 GPa and 3.2%, respectively. Theseresults exceeded the mechanical properties of AS4 unidirectional carbonfiber epoxy composites. The failure strain was double that of AS4composites. After functionalization, the tensile strength and Young'smodulus increased to 3,252 MPa and 279 GPa, respectively. This is 80%and 250% higher than the tensile strength and Young's modulus ofcoagulation-spun, single-walled carbon nanotubes/polyvinyl alcoholcomposite fiber previously reported. The failure strain offunctionalized aligned CNT nanocomposites dropped to 1.6% from 3.2% dueto the chemical bond formation between CNT and epoxy resin. Based onthis measured Young's modulus of aligned and functionalized CNT sheetreinforced epoxy composite, an orientation factor η₀=0.8 can be had,according to the results of Polarized Raman spectra analysis, and a loadtransfer efficiency factor of η_(B)=0.972 as previously discussed.

Hence, Equation (2) can be used to calculate the Young's modulus of CNTbundles. The result was 714 GPa, which is consistent with theexperimental values reported. FIG. 39A shows the fracture surfacemorphology of pristine aligned CNT sheet reinforced epoxy compositespecimen after tensile tests. There are no broken nanotubes observed.FIG. 39B shows the nanotubes separated from the epoxy resin, whichindicates the poor interfacial bonding between pristine CNT and epoxyresin. After functionalization, some of broken nanotubes can be observedat the fracture surface of the functionalized aligned CNT sheet/epoxycomposite, as shown in FIG. 39C, indicating better interfacial bonding.FIG. 39D shows a heavily curved thin film formed of functionalized CNTswell bonded with epoxy resin peeled from the fracture surface, furtherillustrating interfacial bonding improvement. FIG. 39E is the HRTEMimage of cross-section of pristine aligned CNT sheet reinforced epoxycomposites. Most double-walled nanotubes collapsed into “dog-bone” shapeand stacked very well along the alignment direction. The results revealthe intertube frictional force can be increased by a maximum factor of4, when all tubes collapse and the bundle remains collapsed.Furthermore, the bundle will become stronger due to the significantdecreasing of overall cross-sectional area for the collapsed structure.Herein, the collapsed double-walled nanotubes were observed in thepristine aligned CNT sheet reinforced epoxy composite. One reason forcollapse may be the high pressure in the press of fabricating thecomposites. These collapsed nanotubes packed very well, which resultedin high CNT loading and high mechanical properties of CNT sheetreinforced epoxy composites. Normalized to 60% reinforcement volumefraction, the tensile strength of the functionalized and aligned CNTcomposites was 10-20% higher than the state-of-the-art high-strengthunidirectional structural CFRP systems, such as unidirectional T1000Gcomposites, as shown in FIG. 40A, and about 5×, 3× and 2× greater thanthat of aluminum alloys, titanium alloys and steels for structuralapplications, respectively. The Young's modulus of the resultant CNTcomposites was two times higher than typical unidirectional AS4, IM7,T300, T700 and T1000 CFRPs, and close to the best high-modulus CFRPsystems (M55J and M60J graphite fiber composites). The strain of thisnanotube composite was 2 times that of the CFRP systems at the samelevel of Young's modulus, as seen FIG. 40B, which is an improvementtoward developing more resilient composites. The measured density of ourCNT composites was 1.53 g/cm³, slightly less than carbon fibercomposites.

Thus, a new class of resilient, high-mechanical performance nanotubecomposites may be developed by utilizing extra-large aspect ratio CNTs,optimizing alignment and improving interfacial bonding. These compositeswill lead to uncompromised design freedom and unprecedented performanceadvantages for engineered systems in aerospace, automotive, medicaldevices and sporting goods industries. Advantages include weightreduction, high stiffness and strength, great resilience and toughnessfor improved damage tolerance and structural reliability, as well ashigh electrical and thermal conductivity for multifunctionalapplications.

Publications cited herein and the materials for which they are cited arespecifically incorporated by reference. Modifications and variations ofthe methods and devices described herein will be obvious to thoseskilled in the art from the foregoing detailed description. Suchmodifications and variations are intended to come within the scope ofthe appended claims.

We claim:
 1. An article comprising: a sheet comprising a network ofnanoscale fibers and a matrix material, wherein the matrix materialcomprises a resin, wherein the network has been mechanically stretchedto align at least a portion of the nanoscale fibers, and wherein thenetwork has (i) a Young's modulus ranging from about 5 GPa to about 25GPa in the direction of the nanoscale fiber alignment or greater than 25GPa in the direction of the nanoscale fiber alignment, and (ii) atensile strength (a) ranging from about 200 MPa to about 668 MPa in thedirection of the nanoscale fiber alignment, or (b) greater than 668 MPain the direction of the nanoscale fiber alignment.
 2. The article ofclaim 1, further comprising a supporting medium that has beenmechanically stretched with the network.
 3. The article of claim 2,wherein the supporting medium comprises a flexible thermoplasticmaterial.
 4. The article of claim 1, wherein the resin is an epoxyresin.
 5. The article of claim 1, wherein the resin comprises a phenolicresin, a polyimide resin, a bismaleimide resin, a cyanate resin, athermoplastic resin, or a combination thereof.
 6. The article of claim5, wherein the thermoplastic resin comprises nylon, apolyetheretherketone resin, or a combination thereof.
 7. The article ofclaim 1, wherein the network has a tensile strength ranging from about620 MPa to about 2,088 MPa.
 8. The article of claim 1, wherein thenetwork has a tensile strength ranging from about 620 MPa to about 3,081MPa.
 9. The article of claim 1, wherein the network has a tensilestrength greater than 3,081 MPa.
 10. The article of claim 1, wherein thenetwork has a Young's modulus ranging from about 47 GPa to about 169GPa.
 11. The article of claim 1, wherein the network has a Young'smodulus ranging from about 47 GPa to about 350 GPa.